Microstructural characterization of rapidly solidified aluminum transition metal alloys

Materials Science and Engineering, 91 (1987) 201-216

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Microstructural Characterization of Rapidly Solidified Aluminum Transition Metal Alloys J. M. SATER, S. C. JHA and T. H. SANDERS, JR. School of Materials Engineering, Purdue University, West Lafayette, IN 47907 (U.S.A.) (Received July 21, 1986. in revised form September 3,1986)

ABSTRACT

Current research in the field o f rapid solidification is directed toward the development of new and unusual microstructures with improved engineering properties. The most important advancements in the rapid solidification o f aluminum alloys have occurred in the aluminum-transition metal alloy systems and in process development. The low diffusivities of the transition metals in aluminum and the high elastic moduli of the intermetallic phases in these systems result in microstructures having combined properties of high strength, thermal stability and high elastic modulus. Since coarse particle sizes have deleterious effects on ductility and toughness, control of solidification and composition must be exercised to limit the volume fraction of such coarse primary phases. Microstructural formation in the materials produced by chill methods (ribbon, fiber and splat) and spray methods (powder) will be limited by heat transfer: conduction and convection respectively. Other important factors affecting microstructural formation are the thermodynamic and kinetic considerations related to undercooling. Interface stability can be discussed in terms o f theories of constitutional supercooling and morphological stability (Mullins and Sekerka), and the thermodynamics of non-equilibrium solidification for binary alloys in terms of local equilibrium at the interface. Specific microstructural features of both the as-cast and the consolidated and fabricated melt-spun rib bons and powders will be discussed in relation to their respective solidification sequences and processing conditions. Fracture surfaces will be presented since they *Present address: School of Materials Engineering, Georgia Institute of Technology, Atlanta, GA 30332, U.S.A. 0025-5416/87/$3.50

are representative of the underlying microstructural features. Mechanical properties and fracture mechanisms are also discussed in terms of these features. 1. INTRODUCTION The response to demands for reduced weight and increased performance aircraft engines has prompted designers to consider aluminum alloys as replacements for the heavy metals, such as titanium, currently in use. The candidate alloys are required to be stable at elevated temperatures, approaching 500 K, to be most effective as replacement materials. Since conventional high strength 2XXX- and 7XXX-type aluminum alloys cannot be utilized at these temperatures, alloy developers are turning to non-traditional transition metal solute additions, the added quantities being in excess of their solid solubility limits in aluminum. The limited solid solubilities of these elements necessitate the use of novel fabricating approaches. Techniques such as rapid solidification are required to achieve microstructures having acceptable combinations of thermal stability, strength, elastic modulus and fracture toughness for aluminum-transition metal alloys. As early as 1947, investigators showed interest in the development of high modulus, high temperature aluminum alloys. The elements silicon, beryllium, cobalt, manganese and nickel with concentrations in excess of the limits of solid solubility in aluminum were found to offer superior combinations of modulus and strength. The Royal Aircraft Establishment suggested that yield strengths of 21 t o n f in -2 and tensile strengths of 28 t o n f in -2, with a modulus of 12 X 106 lbf in -2 and ductilities of 4% might be possible. However, because of the restrictions on the solidification processes of the day, the size and distri© Elsevier Sequoia/Printed in The Netherlands

202 butions of the intermetallic phases containing these elements were sufficiently coarse to preclude having acceptable ductilities. It was not until the late 1950s that alloys containing large quantities of insoluble elements had the potential to be used as engineering alloys. With the pioneering work at Alcoa Research Laboratories, binary A1-Fe alloys, ternary A1-Fe-Ni and A1-Fe-Cr and quaternary A1-Fe-Cr-V alloys were fabricated into extruded sections from prealloyed atomized powders. Although the microstructures were coarse, they were significantly finer than those produced b y ingot-casting methods. Over the last 25 years, new non-equilibrium alloy phases and aluminum alloy microstructures which exhibit partitionless structures and refined microstructural features have been produced. This paper will focus attention on the influence of concentration and casting parameters on the microstructure of rapidly solidified aluminum alloys and the influence of the as-cast structures on the microstructure and properties of wrought products. Although there has been much work on the microstructures, properties and processing variables of powder metallurgy aluminum alloys based on the 7XXX alloys (aluminum alloys 7090 and 7091) and on the 2XXX alloys, in this paper the microstructure of rapidly solidified aluminum alloys with compositions significantly different from those which can be produced by conventional ingot-casting methods will be emphasized.

2. CHARACTERISTICS OF RAPID SOLIDIFICATION

Solidification from the melt is an essential step in almost every sequence of metal processing. Examples include the production of ingots or continuous strands for working into more useful forms and the manufacture of net-shaped castings for direct application. The properties of the final product are determined to a great extent b y the nature of the solidification process itself. For large castings which undergo conventional solidification processes, the amount of segregation is significant. An advantage, then, of completing solidification in a short (seconds or less) period of time and in thin (millimeters or less) sections is that

there is little time or space within which to form a coarse segregated structure. Heat flow conditions prior to and during solidification have a pronounced effect on achievable interface velocities and, therefore, on the resulting microstructures. Solidification over small dimensions (a small thermal mass) and at high rates can be accomplished b y removing the heat either conductively or convectively. The rapidly solidified forms such as ribbons, fibers and splat are cooled primarily b y conduction. The heat is transferred away by a substrate. The other rapidly solidified form, powder, is cooled primarily by convection. The heat is transferred away from the fine molten droplets in a gas stream. There are similarities and differences in the two thermal histories that will be discussed separately, and the resulting microstructures will be compared.

2.1. Heat transfer in splat and melt-spun materials Both melt-spun and splat rapidly solidified materials experience heat loss b y conduction. The solidification front moves into a superheated liquid, and the heat of fusion is withdrawn through the solid. The achievable interface velocities are limited b y the rate of heat extraction to the substrate and the thickness of the material being processed. The undercooling obtained in melt spinning i s limited, at best, and, because of the positive temperature gradient ahead of the solid-liquid interface, the controlling factor in microstructural development is the high interfacial velocity. One of the problems associated with splat or melt-spinning processes is that a build-up of material on the substrate will reduce the heat transfer and, therefore, the interface velocity. Very coarse microstructures consisting of equilibrium phases may form. Consider the case of a parallel-faced slab of thickness z losing heat from one face, as in a splat rapidly solidified material. The heat lost to the surroundings is governed by the heat transfer coefficient h. If h is low enough to ensure that the temperature remains essentially uniform throughout the slab during cooling and solidification, newtonian conditions will apply. For newtonian conditions, it is assumed that the dominant resistance to heat flow occurs at the surface and that no significant temperature gradients can develop inside

203 the rapidly solidified sample [1]. The heat transfer coefficient h is assumed to be constant. The thermal properties of the solid and the liquid are considered to be different b u t independent of temperature. The newtonian approximation becomes less valid as the sample size and/or h increase(s). The assumF~;on of uniform temperature distribution also becomes less valid when substantial undercoolings prior to nucleation are attained since local recalescence of the interface is faster than the overall recalescence of the sample. At very large undercoolings, heat is being transferred from the solid into the liquid. Utilizing some simple heat transfer equations, Jones [2] found that the cooling rate was directly proportional to h/z, and the solidification front velocity to h. Both the cooling rate and the front velocity are independent of the position x where x is the fraction of z which is solid during solidification. Several researchers [3-5 ] have attempted to analyze the heat and mass transfer and fluid flow conditions existing during melt spinning. Kavesh [6] identified two transport-limiting cases dominating ribbon formation: (a) thermal transport control as described previously and (b) m o m e n t u m transport control. In case (a) the ribbon forms by solidification inside the " p o o l " of liquid located below the orifice on the substrate. For case (b) the liquid film is pulled from the " p o o l " to solidify farther downstream. Because of the large thermal boundary layer width in comparison with the m o m e n t u m or boundary layer thickness, the mechanism of thermal transport has been considered to be of major importance in melt spinning. It has recently been observed that, since h between the melt and the substrate is some finite value, solidification within the " p o o l " is difficult [7 ]. Thus the m o m e n t u m transport mechanism becomes more important. Both mechanisms, however, will affect ribbon formation, i.e. viscosity changes due to temperature reduction within the pool will affect the flow pattern. From results of an analysis involving use of the explicit finite difference method, Takeshita and Shingu [8] concluded the following. (1) For large h (better thermal contact between the substrate and the melt) the ribbon thickness decreases. (2) An increase in substrate velocity decreases the ribbon thickness.

(3) An increase in the initial temperature of the melt may increase the ribbon thickness. Effects vary because of the interaction of thermal and flow transport mechanisms. (4) The cooling rate experienced b y a solidifying ribbon may decrease b y an order of magnitude from the substrate to the freesurface side when h is very large. Thus, ribbon thickness via the external heat transfer conditions will be a critical casting parameter. Clyne [9] has developed a one-dimensional finite difference model using an explicit solution scheme based on truncated Taylor series expansions. He found that values of h, the heat transfer coefficient, range from 103 up to 106 W m -2 K -1 for melt spinning. Increased ribbon thickness (smaller h) increases both solidification time and temperature differences in the system such that it is possible for an initially undercooled liquid to become superheated and, therefore, to solidify more slowly through recalescence. These changes may be accompanied by structural degeneration such as microsegregation or phase separation. To avoid adiabatic recalescence (the net latent heat evolved during recalescence is much greater than the net heat extracted through the surface during recalescence), the Mehrabian number Me must be less than some critical value related to the nucleation temperature (Me is the ratio of the kinetic coefficient which represents the proportionality constant between the crystal growth velocity, undercooling and heat of fusion to the heat transfer coefficient and is given by B z~/f/hi). It is certainly of practical importance to determine the complex effects of h, AT and other process variables to eliminate adiabatic recalescence through improved processing. Considerations of melt-substrate contact time for continuous processes should also enter into such studies.

2.2. Heat transfer in atomized powders Heat transfer conditions inside a supercooled liquid droplet have been studied extensively by Levi and Mehrabian [1, 10]. In the atomization process, heat flow conditions involve solidification into a supercooled melt, where most of the heat of fusion is absorbed by the liquid through recalescence. Achievable interface velocities are limited by the amount of undercooling rather than b y the rate of heat extraction or sample volume.

204 Two distinct solidification regimes can be identified in the atomization process. The first is the rapid solidification stage; the heat of fusion is absorbed by the supercooled droplet. During this stage, heat lost to the environment is negligible. During the second stage the external heat transfer limits the progress of solidification. This stage, one of relatively slow growth, develops after the droplet undercooling has been relieved and is approximately isothermal. Levi and Mehrabian [1, 10] have analyzed the solidification of aluminum droplets using both a newtonian and non-newtonian (enthalpy) model. As mentioned previously, the assumptions of no significant temperature gradients in the droplet and of dominant heat flow resistance at the surface form the basis for the newtonian model. An increase in droplet diameter and substantial undercoolings preceding nucleation decrease the model's validity. Other assumptions include a stable interface configuration throughout solidification and negligible surface tension effects on growth kinetics. Numerical solutions using a one-dimensional finite difference model revealed that, for a given growth geometry, the thermal history of the droplet is independent of particle size but that it can be characterized by Me. For typical Me values in atomization (Me ~ 100), external cooling has a negligible effect in extending the rapid solidification stage to a solid fraction larger than that achieved by the limiting case of adiabatic solidification. An increase in the number of nucleation events also reduces the role of external cooling during recalescence. For the non-newtonian (enthalpy) model, Levi and Mehrabian [1, 10] introduced a nonuniform temperature distribution. The initial temperature profile in the droplet is one of concentric isotherms. The problem is two dimensional because the thermal field during solidification is symmetric about the growth axis. Additional problems of undercooling and solidification rate variations along the interface leading to interface shape deviations may possibly exist. The droplet is divided into volume elements or nodes for a rotational bipolar coordinate system (u, v, w) having coincident rotational and growth axes, and the interface is defined by a series of annular steps propagating along lines of constant ui where u represents a coordinate in the bi-

spherical system perpendicular to the growth direction. The temperature gradient in the solid is restricted to values of zero or more. From the results of the nodal calculations, Levi and Mehrabian [1, 10] concluded the following. (1) Local recalescence allows the interface to change its velocity and its adjacent thermal field to accommodate the heat of fusion. (2) At a given undercooling with increasing particle size the extent of the rapid solidification regime is reduced. (3) Decreasing the particle size and the nucleation temperature and increasing the heat transfer coefficient h will increase the interface velocities and the amount of undercooling. (4) Production of a homogeneous solid requires an undercooling described by the following relationship: Tc

AHm~ f

CldT

(1)

d TN

where AH m is the latent heat of fusion, C1 is the specific heat of the liquid, TN is the nucleation temperature and Te is some critical temperature. As long as the interface temperature is below To, which is less than To [11], the interface remains planar. To obtain a completely homogeneous solid, the droplet must be supercooled to T N o r "hypercooled" such that the heat of recalescence produced never increases the melt temperature above To. The non-newtonian model does appear to offer a more reasonable explanation of the thermal history of a solidifying droplet.

2.3. General thermodynamic and kinetic considerations The stability of the solidification interface associated with rapid growth following undercooling is important in determining the final microstructure. The classical theory of constitutional undercooling has generally been used as a fair approximation for determining growth conditions that result in unstable interfaces [12]. This theory, which mainly describes cases of low front velocities, is essentially an application of thermodynamics to determine whether the liquid ahead of the moving front is supercooled with respect to its concentration. The stability parameter is a ratio of the temperature gradient G1 in the liquid preced-

205 ing the moving front to its velocity R. Increasing the ratio enhances interface stability. A planar interface, necessary for partitionless solidification, is stabilized by a steep temperature gradient and destabilized b y a steep solute gradient resulting from microsegregation of solute and solvent atoms. So, if G1 is negative, any perturbation on the planar interface becomes stable. The product G R is a measure of the average cooling rate during solidification. Increasing values of G R , therefore, correspond to the finer microstructures obtainable in rapid solidification. It should be noted that conventional cellular or dendritic structures are predicted b y this theory for most solidification sequences. For very high interfacial velocities, however, other effects and factors become more important: (1) surface tension and latent heat evolution effects; (2) thermal and diffusional field effects; (3) local equilibrium at the interface. Consideration of these factors led to the development of the theory of morphological stability by Mullins and Sekerka [13, 14] and Sekerka [15 ], which considers interface perturbations with respect to solute and heat diffusion fields and surface tension. The perturbations existing at high interface velocities are so small that capillarity forces acting on the perturbations tend to dissolve them. The planar interface is maintained as long as the net heat flow is toward the solid (KsG s -K1G 1~ 0 where K is the thermal conductivity). If the net heat flow is toward the solid, the only factor destabilizing the planar interface is the solute field ahead of it. This effect, in the regime of high interfacial velocities, is termed absolute morphological stability [4, 5 ]. Another important effect at high front velocities is caused by the departure from local equilibrium at the interface. Coriell and Sekerka [3] concluded that most departures from local equilibrium enhance interface stability, particularly where k -> 1. The partition coefficient k is defined as the ratio of concentration of the solute in the solid to that in the liquid, i.e. CJC1. Baker [16] studied the behavior of k relative to the interface velocity and indicated that k -~ 1 at sufficiently high velocities for most alloys. It has been shown [3, 5] that, at high growth rates and with the growth or decay of a perturbation under consideration, surface energy effects become increasingly important in inter-

face stability compared with the destabilizing effect of the solute gradient. However, there are some shortcomings (which are areas for further research) as follows. (1) The conclusions are based on the use of macroscopic transport theory, which cannot be justified when solidification rates are so rapid that critical diffusion lengths are of the order of atomic dimensions. (2) Results are only strictly applicable to constant velocities, a situation rarely occurring in most experiments. (3) Specific quantitative conclusions cannot be obtained b y this theory when departure from local equilibrium is considered. Morphological stability theory, however, does generally predict far greater stability for rapid solidification than that predicted by constitutional supercooling. One of the most important factors in allowing rapid solidification to occur is the amount of undercooling at the onset of nucleation. In numerous studies b y Perepezko and coworkers [17-20], the amount of liquid undercooling has been found to be a key factor in controlling phase selection kinetics which, in turn, determine microstructural development. A range of possible nucleation sites showing different nucleating capabilities exist in any liquid sample. By continually isolating or removing nucleants from the sample an undercooling level characteristic of the most potent remaining nucleant can be observed [20]. The use of small thin sections or droplets via rapid solidification processing basically allows the number of nucleant sites to be reduced so that large undercoolings can occur. If the amount of undercooling is large enough, and the solidification is fast enough, partitionless solidification or complete solute trapping can occur. Some consideration of thermodynamic conditions is important in any discussion of non-equilibrium solidification microstructures, particularly in terms of partitionless solidification. Baker and Cahn [11] have discussed the thermodynamics of non-equilibrium solidification for binary alloys. The basis of their discussion rests on the assumptions that thermodynamics may be applied locally and that different processes at the interface are independent. From these stems the assumption of constrained equilibrium, i.e. certain processes are assumed to occur or not to occur so slowly that thermodynamics are applicable to the

206 remaining processes. Examples include metastable phase equilibrium (constraint:absence of one or more stable phases) and partitionless solidification (constraint: diffusional mobility of one or more components). With the aid of free energy vs. concentration curves, they defined the domain of possible interface concentration for isothermal and steady state solidification. They concluded that the maximum solid concentration Cs that can be formed from a liquid of concentration CL at a particular temperature is determined b y the point at which the solid and liquid free-energy curves intersect. This is known as the C(To) concentration. The locus of the concentrations and temperatures where the liquid and solid free energies are equal is known as the To curve. It represents the highest interface temperature at which the partition coefficient k = CJCL (ratio of the concentration of solute in the solid to that in the liquid) can be unity. Partitionless transformations, which occur at some T < To for a given concentration, involve an increase in the solute chemical potential on crossing the interface, which, in turn, implies that the species are solidifying in a cooperative manner. Although the overall free energy decreases, a c o m p o n e n t experiences an increase in chemical potential on crossing the interface because (1) it is trapped by the advancing solidification front or (2) it is a necessary participant in an independent solidification reaction mechanism involving several species. Consideration of the reaction kinetics may further decrease the critical temperature below the To value [21]. The major problem with classical thermodynamics in general is its inability to predict precisely what structures will appear. In order to examine the question of local equilibrium at the interface for large growth rates, Baker and Cahn [22] conducted a study of splat-quenched alloys at the zinc-rich end of the Zn-Cd system which exhibits a retrograde solidus (a maximum in the solidus such that C~(eq max) > C~(eq)). Their reasoning for choosing a system with a retrograde solidus is that, if an actual solid concentration could be measured which is greater than the retrograde maximum in the equilibrium solid concentration, a positive departure from local equilibrium at the interface must have occurred. This behavior is what they observed. The question of local equilibrium can be ad-

dressed in part by the use of irreversible thermodynamics, although there is a definite need to study such departures as well as the relationship between the rates of reaction and the thermodynamic quantities b y further experimentation.

3. MICROSTRUCTURAL EFFECTS

3.1. General microstructural effects: rapidly solidified materials Micro structures resulting from rapid solidification generally contain the features of conventional casting including grain, dendritic and eutectic morphologies, although on a much finer scale. The sizes and quantities of microstructural features are strongly dependent on the particle volume through undercooling and interface velocity, as well as on concentration. Thin sections of gun splats may exhibit elongated or equiaxed grains with diameters of a b o u t 0.1 pm [23]. Thicker multilayer regions may exhibit columnar growth extending from layer to layer [24]. Melt-spun ribbons may exhibit an initial chill zone on the substrate side, leading to columnar growth with the possibility of an equiaxed zone adjacent to the air side [25]. Most powder particles, having an average diameter much less than 100 ttm, are single grains, although some A1Si powders have been observed to contain two or more grains and, even, twins [10]. Rapidly solidified materials exhibit extremely fine dendrites, often with no secondary arms, for samples cooled at very high rates. In addition to dendritic and cellular growth, "predendritic" growth has been observed on the chill face of splat and melt-spun rapidly solidified products [26]. These regions are points of contact and nucleation on the chill face from which growth begins with a smooth curved front retaining the original solute content of the parent melt and are surrounded by a region in which the dendritic arms are noncrystallographic. This interface breaks down into a cellular-dendritic mode with the associated solute microsegregation [27 ]. For a totally featureless non-segregated microstructure, partitionless solidification is necessary. Cellular-dendritic structures have also been observed in numerous powder studies.

207 Eutectic microstructures have also been observed in splat and melt-spun materials, as well as in powders. Kofler [28] found that a melt of non-eutectic composition could solidify as a ~quasi-eutectic" over a concentration range that increaseswith increasing undercooling. These regions are referred to as coupled growth zones, the shape of which is determined by growth rate variations with undercooling and concentration [29]. However, because of the requirements o f constant interface velocity and unidirectional heat flow, it is difficult in practice to see a ribbon or a variety of splats having a completely coupled eutectic structure.

3.2. Specific microstructural effects The majority of the product microstructures in rapidly solidified aluminum-transition metal alloys consists of finely divided secondary intermetallic phases. These phases arise from breakdown of the cellular-dendritic as-cast microstructure and from formation of a small fraction of precipitates by decomposition of the supersaturated as-cast microstructure during degassing, preheating and consolidation. This cellular-dendritic breakdown is illustrated in Fig. 1 for an A1-Fe-NiCo alloy. These secondary phases are primarily responsible for the strengthening. A small fraction of the microstructure may contain coarse primary intermetallics which have formed in areas of relatively slow solidification. These coarse particles do not change size or shape significantly on exposure to elevated temperatures, although solute redistribution does occur in the surrounding cellular network. However, the size and spacing of the individual particles formed by solute redistribution are a function of the time at some temperature. The fractions of the various constituents are, thus, a function of the cooling rate. By increasing the cooling rate, the secondary intermetallics can be refined and the coarse primary intermetallics can be eliminated. Optimum tensile properties will result from a rapidly solidified particulate free of coarse primary intermetallics but containing a high volume fraction of fine secondary intermetallics. The solidification structures found in splat and melt-spun aluminum-transition metal alloys have been studied and described in detail by Jones [24] and by Garrett and Sanders [30]. Zone A, consisting of fine primary aluminum dendrites free of optically visible pre-

Fig. 1. Distribution of solute in a relatively slowly cooled A1-3.3wt.%Fe-2.3wt.%Ni-4.6wt.%Co splat (a) before heat treating and (b) after heat treating f o r l h a t 673K.

cipitates, is found on the substrate side of splat and ribbon while zone B, consisting of primary intermetailic particles surrounded by primary aluminum, is generally found on the air side. Primary intermetallics on the air side nucleate simultaneously with primary aluminum on the substrate side. Primary aluminum with solute segregated at dendrite-cell bound. aries surrounds each intermetallic and grows radially outward until reaching other growth fronts. The kinetics of the solidification process suppress nucleation of the intermetallic so that aluminum nucleates and grows. Although solute is found in the intermetallics and segregated in the cellular structure in zone B, and it remains in the supersaturated matrix in zone A, average overall concentrations in each zone would be similar. Primary intermetallic growth is controlled by external heat transfer conditions; thus the ribbon thick-

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ness is critically important with respect to the microstructural features. The faster a material can be cooled and solidified, i.e. the thinner the ribbon or splat, the finer the cells and primary intermetallics will be. Eventually at high enough cooling rates the primary intermetallics are eliminated, and a featureless microstructure is formed. The presence of primary intermetallics can also be attributed to insufficient time above the liquidus temperature which prevents complete dissolution of primary phases. This condition, in some cases, may be responsible for the appearance of primary intermetallics throughout the ribbon thickness. However, in many transmission electron microscopy (TEM) investigations, featureless regions have only been observed on one side of the ribbon, while coarse primary phases have been observed on the air side. If the presence of the intermetallics was caused b y those existing in the melt prior to solidification, they should be randomly distributed throughout the thickness. In the atomized powders, nucleation generally occurs at the droplet surface. A homogeneous region of zone A is produced as solidification proceeds into a highly undercooled melt, provided that the previously mentioned heat conditions are satisfied. The interface velocity and undercooling are reduced as the heat of fusion is absorbed into the liquid. When the interface temperature becomes greater than To, solute rejection into the liquid occurs. Eventually the interface becomes unstable and dendrite-cell growth occurs. Different thicknesses of ribbon containing a variety of solidification structures are produced because of variations in the melt-spinning process. The microstructures observed in the A1-Co system, for example, are clearly affected by. the ribbon thickness and, therefore, b y heat transfer. The microstructures found on the air side of the ribbon are similar for different thicknesses. Since the number of p r i m a r y A I 9 C o 2 intermetallics is large, the grain sizes are small because of limited growth opportunities, although the cell sizes are relatively large. Microstructures found on the substrate side vary from the relatively large primary intermetallics, associated with small grains having coarse cells in a thick ribbon, to smaller primary intermetallics with large grains having finer cells in a thin ribbon. Examples of these coarse structures are shown in Fig. 2.

Fig. 2. Section of Al-10wt.%Co ribbon comparing slowly solidified microstructures. Section (a) was solidified faster than section (b).

Fig. 3. Section of Al-10wt.%Co ribbon showing growth fronts which have undergone the fastest solidification.

Microstructures of the thinnest ribbon consist of growth fronts containing no primary intermetallics. These growth fronts can be seen in Fig. 3. Figure 4 shows a typical through-thick-

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Fig. 4. Typical through-thickness section of A1-Fe-X splat showing the regions that have undergone the highest solidification rates (zone A) adjacent to a zone B region.

ness section o f A1-Fe-X splat showing regions of zone A and zone B. Different powder sizes will also contain a variety of solidification structures. These microstructures are affected by the powder size through the a m o u n t of undercooling. The largest powders or powders containing a high solute c o n t e n t generally show enhanced segregation including, for example, the presence of coarse primary intermetallics and cells in A1-Fe-Ni, A1-Fe-Ni-Co and A1-Mn-Si alloys [31]. Smaller particles may exhibit a totally cellular structure or a combination of cells and zone A. A typical powder structure is shown in Fig. 5. The smallest particles will exhibit a complete solid solution, zone A. The solidification rate is related to undercooling and has an important effect on microstructure: the faster the solidification rate,

the finer is the microstructure. Material undergoing relatively slow solidification contains relatively coarse intermetallics and large cell sizes, while material undergoing higher solidification rates contains progressively smaller intermetallics and cell sizes. At the highest solidification rates (i.e. smallest particle dimensions), partitionless solidification occurs, and no intermetallics or cells are observed. The cobalt in A1-Co ribbons is segregated two distinct ways. For ribbon which has undergone relatively slow solidification the cobait is segregated in the cellular structure surrounding the primary A19C02 intermetallic. On aging, the solute in these regions pinches off to form discrete unoriented A19C02 precipitates. Figure 6 shows the breakdown of the cellular structure in A1-Fe-X powders. In the most rapidly solidified ribbon, cobalt is

210

Fig. 5. Typical powder microstructure of A1-Fe-X showing a possible combination of cellular structure and zone A.

Fig. 6. Typical powder microstructure of A1-Fe-X ~howing the breakdown of the cellular structure and the formation of precipitates at powder boundaries in a consolidated material.

segregated into what appears to be a banded structure radiating f r om the nucleation site. During aging, these segregated bands form oriented plates o f A19Co2 intermetallic (possibly metastable) which resemble a tweed

Fig. 7. Section of Al-10wt.%Co ribbon showing banded cobalt structure in growth fronts which have undergone the fastest sofidification for (a) an as-cast specimen and (b) a specimen aged for 1 h at 573 K.

fabric or Guinier-Preston zones and 0" in A1Cu alloys. T he plate-like precipitates are oriented on (001)-type planes and along [001]t y p e directions of the matrix. This precipitation sequence is shown in Fig. 7. As-cast structures in A1-Fe and A1-Fe-Ce are shown in Fig. 8. J a n o t and Lelay [32] studied the A1-Fe solid solution containing a very small a m o u n t of solute (C < 240 × 10 -6 at.%) using MSssbauer spectroscopy and concluded that a true random solid solution is not obtained in A1-Fe alloys. T h e y also deduced t he presence of solute-rich clusters in the aluminum matrix after annealing their samples at 300 °C. In melt-spun A1-Fe alloys initially containing zone A, spherical solute-rich regions have been observed by TEM following aging [33]. At longer times, these clusters rearrange themselves into an ordered phase which has been identified in A1-4wt.% Fe splat [33, 34] and A1-8wt.% Fe ribbon [33, 35] with the aid of electron diffraction. At even longer aging

211

Fig. 8. (a) Section of A1-8wt.%Fe ribbon showing primary intermetallics and (b) section of A1-8wt.%Fe-8wt.%Ce showing growth fronts which have undergone the fastest solidification (as cast).

times the ordered phase transforms to semicoherent needles of equilibrium A13Fe that eventually lose coherency with the matrix and assume irregular morphology. Addition of cerium to A1-Fe alloys suppresses formation of the metastable ordered precipitate, and coherent needles of A13Fe form directly in the matrix. Ribbons of Al-10 wt.% Co have been consolidated using the Rapid-OmnidirectionalCompaction processing technique developed by Kelsey-Hayes. The processed ribbon has sufficient consolidated strength that TEM foils can be prepared. The deformation temperature is low (533 K), and soaking times are short (less than 1 h). The limited a m o u n t of metal flow does not cause recrystallization. Consequently, it is possible to compare as-cast microstructures found in the plane of the ribbon with the microstructures found through the thickness of the consolidated ribbon. As mentioned previously, a given ribbon may contain zone A or B or both because of process variations. For a ribbon containing only

Fig. 9. Section of compacted rapid omnidirectional compaction (ROC) Al-10wt.%Co ribbon showing the edge of the columnar zone (zone A) to the transition region.

zone B, the intermetallics may have existed prior to melt spinning if, for example, the superheat necessary to dissolve the particles initially is inadequate. Conversely, for a ribbon containing only zone A, critical superheat has been attained. The appearance of a combined zone A-zone B structure seems to discount the notion that zone B existence is only related to superheat and pre-existing nuclei in the melt. In particular, the observation that zone A appears exclusively on the wheel side (in a ribbon containing both zones) precludes the possibility that the sole mechanism of zone B formation is growth of pre-existing nuclei in an insufficiently superheated melt. A through-thickness section or ribbon conraining both zone A and zone B can be seen in Fig. 9. The transition region between zones A and B is small to non-existent. This is due to the rate-limiting competitive growth of the numerous grains on the air side. Coarse intermetallics with small grains do appear near the air side while the columnar regions of zone A appear on the substrate side.

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Most microstructural work on consolidatedfabricated, rapidly solidified aluminum alloys has centered on forms produced from powder and splat [36-38]. Studies b y Sanders e t al. [38], utilizing extruded powder and splat materials, showed that final microstructures are affected b y particulate morphologies and solidification rates as well as by consolidation and extrusion temperatures. Two general but distinct regions, one of large-particle distribution and the other of small-particle distribution, have been observed. The large coarse particles formed during the initial solidification process. These two regions can be observed in Fig. 10. The fabricated powders reveal a more uniform particle size with fewer large particles than the splat product has. This behavior is related to the build-up of splat on the substrate which reduces the heat transfer, thus allowing coarse phases to form. Basic

Fig. 10. Microstructures of sections from (a) air-splat and (b) air-atomized particulates consolidated and

extruded at 675 K.

particle shapes present in these products include (1) large, irregularly shaped particles, (2) small equiaxed particles and (3) small oriented rod-like particles [30, 38]. Increasing the consolidation temperature has the effect of creating in the product microstructure a narrow particle size distribution having fewer small particles. Fractography can be utilized primarily as a screening technique in melt-spun material since the fracture surface accurately reflects the underlying microstructure. By bending or pulling the ribbon, some determination can be made of its ductility. A brittle material is necessary for fragmenting the ribbon, and it is also indicative of the featureless zone A microstructure; the finer the starting structure, the finer is the microstructure after fabrication and the better are the strength and ductility. The fracture surfaces of zone B regions (slow solidification rates) containing the primary AlsCo 2 intermetallics exhibit a ductile dimple appearance, as shown in Fig. l l ( a ) . It is believed that the intermetallics act as stress concentrators, and fracture occurs b y coalescence of microvoids formed at these particles. A combination zone A-zone B ribbon (medium solidification rates) undergoes fracture b y delamination along the columnar grains (brittle) and b y microvoid coalescence at the primary intermetallics. A ribbon containing only zone A (highest solidification rate) undergoes brittle fracture b y delamination along columnar grain boundaries, as shown in Fig. ll(c). Heat treating for 1 h at 573 K has no significant effect on the fracture surface appearance of ribbon which has solidified slowly. The surface appearance is dimpled, and fracture has occurred by microvoid coalescence at intermetalhcs (Fig. 1 l(b)). For material containing a combined zone A-zone B structure the ductile dimple surface shows little change while the surface of the region undergoing brittle fracture begins to take on a more ductile appearance. The surface of material containing only zone A, which undergoes brittle fracture prior to aging, begins to resemble that of a ductile fracture surface (Fig. l l ( d ) ) . This behavior is reasonable; when the precipitation sequence is remembered, microvoid coalescence can occur at the oriented plates, thus giving rise to a more ductile fracture appearance.

213

Fig. 11. Typical ductile fracture surfaces ((a) as cast; (b) after aging for 1 h at 573 K) and typical brittle fracture surfaces ((c) as cast; (d) after aging fo,r 1 h at 573 K) for Al-10wt.%Co ribbon.

Fracture surfaces in the consolidated-fabricated materials are also representative of the underlying microstructures. In A1-Fe-Ni-Co alloys [ 38 ], for example, the fracture surfaces generally exhibit a combination of the following: (1) macroparticles such as large intermetallics and oxides (larger than 25 pm) where fracture often initiates; (2) ductile dimple rupture associated with small and large particles (mean diameter range, 0.1-3 pm); (3) delamination along previous particulate boundaries. A typical ductile dimple rupture and delaminated fracture are shown in Fig. 12. There is a direct relationship between the dimple size and distribution and the particle size and distribution. A similar relationship has been observed for regions undergoing fracture by delamination.

4. MECHANICALPROPERTIES Some picture of interrelationships between processing variables, product microstructures and mechanical properties can be developed

from the diverse group of microstructures studied. Microstructural conditions necessary to attain m a x i m u m strength include a high volume fraction of intermetallics, a fine intermetallic particle size, a uniform particle distribution and a thermally stable particle. Thus, fabricated products having a high volume fraction of fine, uniformly sized particles should exhibit the highest yield and tensile strengths. This result has been confirmed by Sanders e t al. [38] and Paris e t al. [31] and is summarized in Table 1. Microstructures of extrusions fabricated from air-atomized powder were found to have the smallest, most uniform particles. These products also exhibited the highest strengths. Paris e t al, [31] have shown that increasing the solute level from 5.0 to 7.5 at.% for the same processing conditions increases the volume fraction of particles. A simultaneous reduction in interparticle spacing occurs, a result which has the effect of increasing the yield and tensile strengths over those of the lower solute alloys. Because of the possible applications of these rapidly solidified wrought alloys in the aero-

214

space industry, some indication of the toughness must be developed to ensure a usable product. The notched-tensile-strength-to-yieldstrength ratio NTS/YS is one such indicator. Generally speaking, a decrease in NTS/YS occurs with an increase in yield strength. The relationship of NTS/YS vs. YS has been found to be similar for the materials studied by Paris e t al. [31] and Sanders e t a l . [38] and is shown in Fig. 13. Alloys containing up to 25 vol.% second phase show an increased strengthtoughness relationship, while the reverse is observed for alloys containing more than 25 vol.% second phase. Factors which could reduce fracture toughness include (1) the existence of uniformly distributed macroparticles and (2) mechanically weak particulate boundaries along which delamination may occur. Another factor m a y be that, as the volume fraction of second phase is increased, the interparticle spacing must decrease, eventually to the point at which plastic deformation cannot be accommodated in the ductile matrix [39].

5. F R A C T U R E M E C H A N I S M S Fig. 12. F r a c t o g r a p h s s h o w i n g regions o f (a) ductile d i m p l e r u p t u r e for an argon-splat A 1 - 3 . 3 w t . % F e 2 . 3 w t . % N i - 4 . 6 w t . % C o particulate c o n s o l i d a t e d a n d e x t r u d e d a t 675 K and (b) d e l a m i n a t i o n for an airsplat A 1 - 3 . 3 w t . % F e - 2 . 3 w t . % N i - 4 . 6 w t . % Co particulate c o n s o l i d a t e d and e x t r u d e d at 675 K.

The effects of macroparticles and particle "bands" on fracture characteristics of A1-FeNi-Co alloys have been studied by pulling tensile specimens without fracturing them

TABLE 1 C o m p o s i t i o n s o f p o w d e r metallurgy alloys

Mn

Si

Fe

Ni

Co

(at.%)

(at.%)

(at.%)

(at.%)

(at.%)

.

. 1.67 1.67 1.67 .

Total solute

Volume fraction

(at .% )

(vol.%)

1.67 2.22 1.11

5.00 5.00 5.00 5.00 6.00

0.41 a 0.24 b 0.24 0.24 0.49 0.29

5.00 ---6.00

. ---.

--

--

2.00

2.00

2.00

6.00

--

--

3.00

3.00

--

6.00

0.29 c

7.50

.

7.50

0.61

.

.

. 1.67 1.11 2.22 .

.

.

5.00

2.50

--

--

--

7.50

0.24 d

--

--

2.50

2.50

2.50

7.50

0.36

6.00

3.00

--

--

--

9.00

0.29

aCalculated bCalculated cCalculated dCalculated

assuming 25 vol.% AII2Mn and 75 vol.% Al6Mn. for Al9(Fe , Ni, Co)2. assuming Al9FeNi. assuming Al12Mn3Si.

215 i

I

"'. 2.0 u) >.

i

e2OVol.

-e~° ~ ' ~ _ ~ j

%

o24

1.5

Iz 1.0

• 2 9 VoL % " - . 0.5

o 36

n'~--.Q "

"~

A41

a4~)

^

"-.

i

J

I

250

350

450

YS

o x i d e characteristics and t h e t h e r m o m e c h a n ical processing o f the m a t e r i a l . Challenges in t h e area o f r a p i d solidification r e m a i n significant. T h e yield o f finer, r a p i d l y solidified p a r t i c u l a t e having m o r e c o n s i s t e n t as-cast m e t a l l u r g i c a l s t r u c t u r e s m u s t be m a x imized. C o n s o l i d a t i o n costs m u s t be lowered. Alloy compositions impractical for convent i o n a l casting t e c h n i q u e s , such as the alum i n u m - t r a n s i t i o n m e t a l alloys, n e e d to be c o n t i n u a l l y identified a n d e x p l o i t e d . Additionally, a b e t t e r q u a n t i t a t i v e u n d e r s t a n d i n g o f s t r u c t u r e - p r o p e r t y r e l a t i o n s h i p s needs to be p u r s u e d so t h a t rapid solidification techniques can be used t o t h e i r fullest advantage.

(MPa)

REFERENCES Fig. 13. Longitudinal NTS/YS vs. YS plot for the alloy extrusions noted in Table 1 (powder and splat particulate).

into t h e region o f plastic i n s t a b i l i t y [ 38 ]. Fine b a n d s o f shear have b e e n o b s e r v e d in t h e small-particle regions; coarse b a n d s o f shear have b e e n seen in the large-particle regions and at stress c o n c e n t r a t o r s . Closer o b s e r v a t i o n of t h e s e surfaces reveals t h e e x i s t e n c e o f microvoids. E x t r e m e m i c r o v o i d f o r m a t i o n a n d c o a l e s c e n c e o c c u r in t h e large-particle regions. T h e s e m i c r o v o i d s initiate at i n t e r f a c e s b e t w e e n t h e large p r i m a r y i n t e r m e t a l l i c particles and coalesce during processing. T h e resulting separated particles f u r t h e r c o n c e n t r a t e stress in these large-particle regions. T h e small-particle regions, h o w e v e r , s h o w n o o b v i o u s m i c r o v o i d formation. T h e f r a c t u r e p a t h m a y involve t h e following. (1) Large void f o r m a t i o n at m a c r o p a r t i c l e s causes strain localization. T h e voids r e d u c e t h e cross-sectional area c a r r y i n g t h e applied load and lead to p r e m a t u r e m e c h a n i c a l failure. (2) Strain l o c a l i z a t i o n occurs in t h e plastically d e f o r m e d coarse particle regions. A t s o m e p o i n t d u r i n g f u r t h e r plastic d e f o r m a t i o n , voids f o r m in these regions at t h e p a r t i c l e - m a t r i x boundary. Delamination along weak boundaries c o n n e c t s t h e s e plastically d e f o r m e d regions. (3) T h e f r e q u e n c y of d e l a m i n a t i o n , contrib u t i n g to strain c o n c e n t r a t i o n and failure, will d e p e n d on t h e s t r e n g t h or w e a k n e s s o f the int e r p a r t i c u l a t e b o n d as related to t h e surface

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